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Nanocomposites containing a thermoplastic blend and organophilic layered clay (organoclay) were produced by melt compounding. The blend composition was kept constant [polyamide 6 (PA6) 70 wt % ϩ polypropylene (PP) 30 wt %], whereas the organoclay content was varied between 0 and 10 wt %. The mechanical properties of the nanocomposites were determined on injection-molded specimens in both tensile and flexural loading. Highest strength values were observed at an organoclay content of 4 wt % for the blends. The flexural strength was superior to the tensile one, which was traced to the effect of the molding-induced skin-core structure. Increasing organoclay amount resulted in severe material embrittlement reflected in a drop of both strength and strain values. The morphology of the nanocomposites was studied by scanning electron microscopy (SEM), transmission electron microscopy (TEM), energy-dispersion X-ray analysis (EDX), and X-ray diffraction (XRD). It was established that the organoclay is well dispersed (exfoliated) and preferentially embedded in the PA6 phase. Further, the exfoliation degree of the organoclay decreased with increasing organoclay content.

Summary: Hydrogenated acrylonitrile butadiene rubber (HNBR) was melt compounded with montmorillonite (MMT) and organophilic modified MMTs prior to sulfur curing. In contrast to the micro‐composite formation resulting from the compounding of the HNBR and pristine MMT, the modified MMTs (i.e., octadecylamine: MMT‐ODA, octadecyltrimethylamine: MMT‐ODTMA, methyltallow‐bis(2‐hydroxyethyl) quaternary ammonium: MMT‐MTH intercalants) produced nanocomposites. It was found that the organoclay with primary amine intercalant (cf. MMT‐ODA) gave confined structures along with the exfoliated/intercalated structures. This was traced to its reactivity with the curatives. By contrast, the organoclays containing less reactive quaternary ammonium compounds (cf. MMT‐ODTMA, MMT‐MTH) were exfoliated and intercalated based on X‐ray diffraction (XRD) and transmission electron microscopy (TEM) results. The hydroxyl functional groups of the MMT‐MTH supported the clay dispersion. The better adhesion between MMT‐MTH and HNBR was explained by hydrogen bonding between the hydroxyl groups of the intercalant and the acrylonitrile group of the HNBR matrix. This HNBR/MMT‐MTH nanocomposite showed the best mechanical properties as verified by tensile mechanical tests and dynamic mechanical thermal analysis (DMTA). The high tensile strength along with the high elongation at break for the rubber nanocomposites were attributed to the ability of the ‘clay network’ to dissipate the input energy upon uniaxial loading.

Natural rubber (NR), polyurethane rubber (PUR), and NR/PUR-based nanocomposites were produced from the related latices by adding a pristine synthetic layered silicate (LS; sodium fluorohectorite) in 10 parts per hundred parts rubber (phr). The dispersion of the LS latices in the composite was studied by X-ray diffraction (XRD) and transmission electron microscopy (TEM). Further information on the rubber/LS interaction was received from Fourier transform infrared spectroscopy (FTIR) and dynamic mechanical thermal analysis (DMTA). Tensile and tear tests were used to characterize the performance of the rubber nanocomposites. It was found that LS is more compatible and thus better intercalated by PUR than by NR. Further, LS was preferably located in the PUR phase in the blends, which exhibited excellent mechanical properties despite the incompatibility between NR and PUR. Nano-reinforcement was best reflected in stiffness-and strength-related properties of the rubber composites.

The deformation behavior of poly(ether ester) is studied by means of small- and wide-angle X-ray scattering (SAXS and WAXS). The material under investigation represents a polyblock thermoplastic elastomer of poly(ether ester) (PEE) type. It comprises poly(butylene terephthalate) (PBT) as hard segments and polyethylene glycol (PEG) as soft segments in a ratio of 57/43 wt %. Isotropic PEE bristles are drawn to five times of their initial length and subsequently annealed with fixed ends for 6 h at 170 °C in vacuum. The WAXS patterns were registered by a pinhole camera and a 2D area gas detector. These measurements were performed both under stress and during the subsequent relaxation in the absence of stress. The deformation was increased stepwise up to the breaking point of the sample (ca. 185%). SAXS patterns were obtained in the same deformation range by means of monochromatic X-ray radiation in the beamline A2 of the synchrotron DESY in Hamburg, Germany. SAXS patterns were registered by means of a 2D “Image-plate” detector. Five deformation intervals were revealed by SAXS. In the first one (ε = 0−50%) an ensemble of uncorrelated strained microfibrils exists and the corresponding layer line small-angle pattern is observed. These microfibrils scatter independently. In the second interval (ε = 50−80%) interactions between neighboring microfibrils develop, a microfibrillar network is observed, and the layer line pattern transforms into a four-point diagram. In the third interval (ε = 80−100%) two additional reflections show up and an unique six-point pattern is seen. Pull-out of tie molecules from crystallites begins to fibrillate the network. This pull-out mechanism is independently proved by WAXS. In the fourth interval (ε = 100−130%) the mean long period of the four-point pattern decreases and the pattern itself vanishes. At last the fiber is completely fibrillated and only a two-point pattern remains visible. In the second, third, and fourth intervals, the microfibrils correlate in the transverse direction, which allows determination of the interfibrillar distance (so-called transverse long period). It decreases gradually with the progress of deformation in both the strained and relaxed state. In the fifth interval (ε > 130%) the long period of the two point pattern remains constant, but its intensity decreases until the fiber breaks. Only a few microfibrils are simultaneously carrying the load. They are destroyed one by one, until the fiber breaks as a whole.

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